Fe-ni-mn-al-cr alloys and methods for production thereof

ABSTRACT

Alloys including iron, nickel, manganese, aluminum and chromium are disclosed. The alloys have high strength and ductility. The alloys are prepared from readily available transition metals, and can be used in applications where properties similar to steel are necessary or advantageous.

RELATED APPLICATIONS

This application is a continuation in part application of U.S. application Ser. No. 12/867,712, international filing date Feb. 13, 2009 which is hereby incorporated by reference and which is the national stage of International Application No. PCT/US09/34123, filed Feb. 13, 2009, which claims the benefit of priority to U.S. Provisional Patent Application Ser. No. 61/028,809, filed Feb. 14, 2008, which is incorporated by reference herein.

GOVERNMENT INTERESTS

The United States Government has rights in this invention under Contract No. NSF-DMR-0505774 and NSF DMR-0905229 between the National Science Foundation (NSF) and Dartmouth College and also under Contract DE-FG02-07ER46392, between the U.S. Department of Energy and Dartmouth College.

BACKGROUND

1. Field of the Invention

This invention relates to novel alloys and methods of producing the alloys. More specifically, the alloys are strong and ductile microstructured alloys having lamellar structures.

2. Description of the Related Art

Basic research on alloy materials seeks to find improved materials, such as those that are lighter, stronger or less expensive than conventional metals and alloys. In other contexts, improved materials may have increased resistance to weather, chemicals or friction in an intended environment of use. Equipment that incorporates these new materials in component parts may have a longer service life, require less maintenance or achieve an improved performance level. From a cost of manufacture standpoint, it is desirable for these new materials to be made from readily available and highly affordable natural resources.

One technique that may be used to produce an alloy with enhanced strength and ductility is a eutectic transformation. A eutectic transformation occurs when components of an alloy crystallize simultaneously from a liquid solution. Products of a eutectic transformation can often be identified by their lamellar structure where spacing between lamellae is typically on the order of less than a micron to a few microns. Such structures are generally strong and ductile. For example, the most well known lamellar material is carbon steel.

SUMMARY

Alloys of the present disclosure advance the art by providing materials with exceptional strength and ductility.

In one embodiment, an intermetallic composition formed by a eutectic transformation in at least two distinct structural phases has an average composition comprising from 25% to 35% iron, 15% to 25% nickel, 30% to 40% manganese and 10% to 20% aluminum, where the composition is described in terms of atomic percentages.

In an embodiment, the invention provides an alloy comprising at least two distinct structural phases, wherein the average composition of the alloy comprises from 25% to 35% iron, from 15% to 25% nickel, from 30% to 40% manganese, 10% to 20% aluminum and the composition is described in terms of atomic percentages. In an embodiment, the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum. In an embodiment, one of the distinct structural phases is a face-centered cubic (f.c.c.) phase. In an embodiment, the concentration of iron and of manganese in the f.c.c. phase is greater than the amount of nickel or of aluminum. In an embodiment, another of the distinct structural phases is a body-centered cubic (b.c.c.) phase, such as B2, an ordered phase. In an embodiment, the concentration of nickel and of aluminum in the B2 phase is greater than the amount of iron or of manganese. In an embodiment, at least a portion of the B2 phase is present in the form of rod-like or plate-like structures. In embodiments, the characteristic width or thickness of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the B2 phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm. In an embodiment, at least a portion of the B2 and f.c.c. phases are present in the form of lamellar structures. Typically at least some of the B2 lamellae alternate with the f.c.c. lamellae as is characteristic of eutectic structures. The characteristic thickness of the f.c.c. phase can be from 200 nm to 2000 nm or from 300 nm to 1000 nm.

In one embodiment, an intermetallic composition formed by a eutectic transformation in at least two distinct structural phases has an average composition according to the formula:

Fe_(a)Ni_(b)Mn_(c)Al_(d)M_(e),

where (in atomic percent) a ranges from 25 to 35; b ranges from 15 to 25; c ranges from 30 to 40, d ranges from 10 to 20, e ranges from 0 to 5 and M is selected from the group consisting of Cr, Mo, C and combinations thereof.

In an aspect, the invention provides an alloy comprising chromium in addition to Fe, Ni, Mn and Al. In an embodiment, the invention provides an alloy comprising at least two distinct structural phases, wherein the average composition of the alloy comprises from 25% to 35% iron, from 15% to 25% nickel, from 30% to 40% manganese, 10% to 20% aluminum and from greater than 0 to less than or equal to 5% Cr; from 2.5% to 5% Cr, from 3% to 5% Cr, from 4% to 5% Cr, from 4% to less than 8% Cr, from greater than 5% to less than 8% Cr, from greater than 5% to 7.5% Cr, from greater than 5% to less than or equal to 7% Cr, from 5.5% to 7.5% Cr, or from 5.5% to 6.5% Cr. In an embodiment, the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum in addition to chromium. In an embodiment, one of the distinct structural phases is a face-centered cubic (f.c.c.) phase. In an embodiment, the concentration of iron and manganese in the f.c.c. phase is greater than the amount of nickel or aluminum. In an embodiment, another of the distinct structural phases is a body-centered cubic (b.c.c.) phase, such as B2, an ordered phase. In an embodiment, the concentration of nickel and aluminum in the B2 phase is greater than the amount or iron or manganese. In different embodiments, greater than 50 at %, from 75% to 100%, from 80% to 95%, or from 80% to 90% of the chromium is located in the f.c.c. phase.

In an embodiment, at least a portion of the B2 phase in the chromium-containing alloys is present in the form of rod-like or plate-like structures. In embodiments, the characteristic width or thickness of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the B2 phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm. In an embodiment, at least a portion of the B2 and f.c.c. phases are present in the form of lamellar structures. Typically at least some of the B2 lamellae alternate with the f.c.c. lamellae. In embodiments, the characteristic thickness of the f.c.c. phase can be from 200 nm to 2000 nm or from 300 nm to 1000 nm. It has been found that the structure of alloys at higher chromium concentrations can differ from that at lower concentrations, with this change in structure being associated with decreased ductility. For example, Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 8 at % Cr produced a finer more complicated structure including cuboidal particles (see Example 4 and FIG. 15B).

In an aspect of the invention, the addition of chromium to the alloy results in improved resistance to hydrogen embrittlement in air at relatively low strain rates. In embodiments, the elongation to fracture of the chromium-containing alloy may be greater than 10% or from 10% to 25% as measured in air at room temperature at a strain rate of 5×10⁻⁴ s^(−1.) These elongation to fracture values may be obtained for dog-bone shaped samples with a gauge length of 10 mm. The ductility of the alloy may also be indicated by the fracture mode of a specimen of the alloy. In an embodiment, the fracture surface indicates ductile tearing with elongated dimples when the alloy specimen is tested in air at room temperature at a strain rate of 5×10⁻⁶ s⁻¹. In an embodiment, the atomic percentage of chromium is from 4% to less than 8% Cr, from greater than 5% to less than 8% Cr, from greater than 5% to less than or equal to 7% Cr, from 5.5% to 7.5% Cr, or from 5.5% to 6.5% Cr.

In an embodiment, the addition of chromium to the alloy may result in some softening of the alloy, but acceptable values of yield strength, hardness and ultimate tensile strength (UTS) of the alloy may be retained. In embodiments, the yield stress of the alloy may be from 600 to 800 MPa or 700 to 800 MPa. The UTS may be from 900 to 1000 MPa. The average hardness of the alloy, as measured by the Vickers Pyramid Number (VPN), may be from 200 to 325 kg/mm².

In one embodiment, a method of producing an intermetallic composition includes heating a mixture of metals, to create a homogenous solution, according to the formula:

Fe_(a)Ni_(b)Mn_(c)Al_(d)M_(e),

where (in atomic percent) a ranges from 25 to 35; b ranges from 15 to 25; c ranges from 30 to 40; d ranges from 10 to 20; e ranges from greater than 0 to less than 8; and M is selected from the group consisting of Cr, Mo, C and combinations thereof; cooling the homogenous solution to obtain a homogeneous solid; reheating the solid to a eutectic transformation temperature; and holding the eutectic transformation temperature for a period of time.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is an X-ray diffraction pattern of Fe₃₀Ni₂₀Mn₃₅Al₁₅ as cast, according to an embodiment.

FIG. 2 is a differential thermal analysis (DTA) curve showing thermal processes of Fe₃₀Ni₂₀Mn₃₅Al₁₅.

FIG. 3 shows a plot of yield strength versus temperature for Fe₃₀Ni₂₀Mn₃₅Al₁₅.

FIG. 4 shows a stress versus strain curve for Fe₃₀Ni₂₀Mn₃₅Al₁₅.

FIG. 5 shows a comparison of tensile and yield strengths for Fe₃₀Ni₂₀Mn₃₅Al₁₅ versus some known alloys.

FIG. 6 shows a comparison of percent elongation for Fe₃₀Ni₂₀Mn₃₅Al₁₅ versus some known alloys.

FIGS. 7A-B illustrate the microstructure of an alloy with the nominal composition Fe₃₀Ni₂₀Mn₃₅Al₁₅. FIGS. 7A and 7B show scanning electron microscope (SEM) images.

FIG. 8 shows a transmission electron microscopy (TEM) image of a quenched Fe₃₀Ni₂₀Mn₃₅Al₁₅ alloy.

FIGS. 9A-C. SEM images of (a) Fe₂₈Ni₁₈Mn₃₃Al₂₁; (b) Fe₂₉Ni₁₉Mn₃₄Al₁₈ and (c) Fe₂₈Ni₂₁Mn₃₃Al_(18.)

FIGS. 10A-D show SEM images of: (a) Fe₃₁Ni₁₈Mn₃₈Al₁₃; (b) Fe₂₉Ni₁₉Mn₃₈Al₁₄; (c) Fe₃₃Ni₁₉Mn₃₄Al₁₄ and (d) Fe₃₆Ni₁₈Mn₃₃Al_(13.)

FIG. 11 shows stress-strain curves of FeNiMnAl alloys with compositions as indicated.

FIG. 12A-D: SEM images of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 2, 4, 6 and 8 at. % Cr (A-D respectively).

FIG. 13A-B. XRD pattern of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with (a) 6 at. % Cr, and (b) 8 at. % Cr, showing peaks for B2 and f.c.c. phases.

FIGS. 14A-E: (a) Bright Field (BF) TEM image of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 4 at. % Cr; (b and c) SAD patterns from the two phases, showing that they are B2 and f.c.c. (the 100 superlattice diffraction spot in the B2 pattern indicates its ordered nature); (d and e) X-ray spectra from the f.c.c. and B2 phases respectively, showing that the Cr resides mainly in the f.c.c. phase.

FIGS. 15A-B: BF TEM image of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 8 at. % Cr: (a) showing two-phase microstructure, and (b) the cuboidal particles accompanied by corresponding SAD pattern, showing that they are B2.

FIGS. 16A-B: (a) Strain-stress curves for Fe₃₀Ni₂₀Mn₃₅Al₁₅ containing different amounts of Cr tensile tested at a strain rate of 5×10⁻⁴ s⁻¹ at room temperature, and (b) UTS, yield stress and elongation as a function of atomic % Cr added to Fe₃₀Ni₂₀Mn₃₅Al₁₅.

FIGS. 17A-B: Strain-stress curves for Fe₃₀Ni₂₀Mn₃₅Al₁₅ containing 6 at. % Cr, tensile tested at different strain rates at room temperature, and (b) elongation to fracture, yield stress and ultimate tensile stress of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr as a function of strain rate.

FIGS. 18A-D: SEM images of fracture surface after tensile testing at a strain rate of 5×10⁻⁶ s⁻¹ at room temperature for: 6% Cr modified Fe₃₀Ni₂₀Mn₃₅Al₁₅ (FIGS. 18A-B) and Cr free Fe₃₀Ni₂₀Mn₃₅Al₁₅ (FIGS. 18C-D).

FIGS. 19A-B: (a) Strain-stress curves of as-cast Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr tensile tested at variety of temperatures at a strain rate of 5×10⁻⁴ s⁻¹, and (b) plot of yield stress and ultimate tensile strength as a function of temperature. The yield stress of Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅ is shown for comparison.

FIG. 20 Elongation to fracture and reduction in area at the neck for as-cast Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr versus temperature.

DETAILED DESCRIPTION

Alloys that are both strong and ductile at room temperature are disclosed, along with methods for making the alloys by way of a eutectic transformation. In some embodiments, observed tensile and yield strengths of the alloys are greater than those of typical stainless steels, and the alloys show ductility comparable to high-strength ferritic stainless steels.

The terms “alloy”, “intermetallic compound” and “intermetallic composition” are used interchangeably herein. They refer to compounds containing at least two different metals.

A “eutectic alloy” is an alloy that is formed when at least two different metals, as well as any non-metals, are present in suitable concentrations and held at a eutectic transformation temperature for a suitable period of time. At the transformation temperature, at least two phases can simultaneously crystallize from a liquid solution to form lamellae of the two phases.

“Room temperature” refers to 20-25° C. as used herein.

The alloys disclosed herein may be used in the manufacture of machine, building and industrial parts. The alloys may be particularly suitable for applications requiring high-strength, wear resistant parts including but not limited to: engines, bearings, bushings, stators, washers, seals, rotors, fasteners, stamping plates, dies, valves, punches, automobile parts, aircraft parts, building materials, and drilling and mining parts. Further, the alloys can be used in any known application currently utilizing stainless steel or any high-strength, ductile alloy.

In a particular embodiment, an alloy contains iron, nickel, manganese and aluminum to which may be added chromium, molybdenum, carbon and combinations thereof. Such an alloy is represented by a macroscopic average formula:

Fe_(a)Ni_(b)Mn_(c)Al_(d)M_(e),  Formula (1)

where M is an alloying addition of any element or combination of elements;

a ranges from 25 to 35;

b ranges from 15 to 25;

c ranges from 30 to 40;

d ranges from 10 to 20;

e ranges from 0 to 5; and

where a-e are expressed on an atomic percent basis.

In one aspect, M may be a metal or combination of metals. For example, M may be chromium, molybdenum, carbon and combinations thereof. In some embodiments, the portion of the alloy that is allocated to M may also range from 0.05 to 4% or in other aspects from 0.5% to 3%.

A narrower formulation that is within the general scope of Formula (1) is:

Fe_(x)Ni_(50-x)Mn_(50-y)Al_(y),  Formula (2)

wherein x ranges from 25 to 35 (atomic percent basis) and y ranges from 10 to 20 (atomic percent basis).

In another aspect, the composition of Fe_(a)Ni_(b)Mn_(c)Al_(d)M_(e) may be within the ranges:

a ranges from 27 to 33;

b ranges from 17 to 23;

c ranges from 32 to 38;

d ranges from 12 to 18; and

e ranges from 0 to 2.5;

where a-e are expressed on an atomic percent basis and M is an alloying addition of any element or combination of elements.

The alloy may be formed by a heat treatment process that results in a eutectic transformation leaving at least two intermetallic phases of different structure and stoichiometry. The macroscopic formulas above pertain to the overall composition, but the macroscopic composition has nanostructure or microstructure of localized phase variances in composition and ordering. The presence of two phases present as lamellae results in ductility along the planes of the lamellae. This ductility may be measured as percent elongation.

In an embodiment, at least one of the phases present in the alloy is in the form of an elongated or lamellar structure. In different embodiments, the phase may be rod-like, plate-like, or a combination thereof. As used herein, a rod-like structure need not be perfectly circular in cross-section, but has a width less than its length. In different embodiments, the elongated structure may have a characteristic length which is greater than its characteristic width, greater than twice its characteristic width, greater than five times its characteristic width, or greater than ten times its characteristic width. As used herein, a lamellar or plate-like structure has a thickness less than its width or length. A characteristic dimension of a structure may be determined from image analysis of a polished section of the alloy. For example, the characteristic width or thickness of a first phase may be determined by measuring the linear distance (within the first phase) between intersections of that phase with a second phase. The distance between structures of the first phase may be determined by measuring the linear distance (outside the first phase) between intersections of the first phase with a second phase.

The rod-like or plate-like structures need not be perfectly straight. However, in an embodiment, adjacent regions of a structure do not form an angle of 90 degrees or less with respect to each other. Structures forming an angle of 90 degrees or less are visible in FIG. 8, where some of the B2 and f.c.c. structures are sufficiently curved that they enclose one another.

The shape and size of the phases in the alloy can influence the mechanical properties of the alloy. If the B2 phase is rod-like in form, the alloy may be viewed as comprising B2 rod-like structures in a matrix of the f.c.c. phase. If both the B2 and f.c.c. phases are plate-like in form, the alloy may be viewed as comprising alternating lamellae of B2 and f.c.c. In an embodiment, the chromium containing alloys of the invention may comprise both plate-like and rod-like B2 structures. In an embodiment, the volume fraction of the plate-like structures is greater than the rod-like structures. In an embodiment, a plurality of the rod-like structures are aligned with one another when viewed in a section approximately parallel to their longitudinal axes. In an embodiment, a plurality of the plate-like structures are aligned with one another when viewed in a section approximately transverse to the plate thickness. In different embodiments, at least 5 or at least 10 adjacent structures may be aligned with each other along at least a portion of the structure. In an embodiment, elongated structures are aligned with one another when their corresponding axes are aligned within ten degrees. In another embodiment, the length of the elongated structures may be greater than 1 micrometer, greater than 2 micrometers or greater than 5 micrometers. In embodiments, the average width or thickness structures of the B2 phase may be from 150 nm to 750 nm or from 150 nm to 500 nm. In embodiments, the average spacing between the structures of the f.c.c. phase can be from 400 nm to 2000 nm or from 450 nm to 1000 nm.

The following examples set forth preferred materials and methods for use in making the disclosed alloys. These examples teach by way of illustration, not by limitation, and should not be interpreted as unduly narrow.

Example 1 Preparation and Characterization of Fe₃₀Ni₂₀Mn₃₅Al₁₅

A quaternary alloy of Fe₃₀Ni₂₀Mn₃₅Al₁₅ composition was prepared by well known arc melting and casting techniques. A quantity of material including 24 g Fe, 17 g Ni, 27 g Mn and 5 g Al was placed in a water-cooled copper mold and heated until molten using the arc melting technique. Ingots were flipped and melted a minimum of three times under argon to ensure mixing. Quenching was done by allowing the alloy to rapidly cool in the copper mold to a temperature of ˜30° C. in approximately 10 minutes. A eutectic transformation was carried out by holding the quenched ingots at about 1215° C. for about 30 minutes. For this composition, the eutectic transformation temperature, as shown in the differential thermal analysis curve, was between about 1210-1290° C., or between about 1212-1250° C. or between about 1214-1230° C. In some embodiments, a 5% excess of Mn may be added to the starting materials because Mn accounts for the majority of weight loss during casting, which results from brittle sharding and evaporation.

The resulting alloy had microstructure in the form of lamellae formed as two intermetallic phases. One phase was a B2 (ordered body-centered cubic, b.c.c.) phase having a composition of Fe₇Ni₄₇Mn₁₈Al₂₈ in terms of atomic percent. The other phase was a face-centered cubic (f.c.c.) phase having a composition of Fe₅₀Ni₇Mn₃₇Al₆ in terms of atomic percent. The widths of the body-centered cubic and face-centered cubic phases were 200 nm and 500 nm, respectively.

The alloy was characterized using analytical techniques that are well known in the art. For example, water displacement analysis was used to determine that the alloy had a density of about 7.02 g/cm³, and chemical composition was determined by energy dispersive spectroscopy (EDS). As discussed above, the overall composition, Fe₃₀Ni₂₀Mn₃₅Al₁₅, was based on microstructured phases of Fe₇Ni₄₇Mn₁₈Al₂₈ and Fe₅₀Ni₇Mn₃₇Al₆.

Structural data was obtained using a Siemens D5000 X-ray Diffractometer with a Kevex silicon detector in the range of 20-110° 2θ, using an instrument that was calibrated against an alumina standard purchased from the National Institute of Standards (NIST). FIGS. 1 and 2 show X-ray diffraction patterns for as cast and quenched samples, respectively. The as cast sample has undergone a eutectic transformation, and displays peaks representative of both the B2 and f.c.c. phases, as shown in FIG. 1.

Room temperature hardness of the two phase alloy, Fe₃₀Ni₂₀Mn₃₅Al₁₅, was determined by taking the average of five measurements from a Leitz Microhardness Indentor with a 200 g load. The average Vicker's hardness was 310 kg/mm².

Differential thermal analysis (DTA) was performed on a Perkin Elmer Pyris Diamond TG/DTA. A typical DTA curve is shown in FIG. 2. The differential thermal analysis suggests that a pre-existing B2 phase begins to disorder at about 1160° C. and a eutectic transformation begins at about 1215° C. This eutectic transformation forms the B2 (b.c.c.) and f.c.c. phases of the lamellar alloy.

Transmission electron microscopy (TEM), performed on either a JEOL 2000FX or a Philips CM 200, indicated that, after formation, the eutectic microstructure was stable up to at least 760° C. Optical microscopy confirmed that no distinct differences were observed between samples that had been annealed for 30 minutes at temperatures between 327-727° C. and then subsequently quenched.

Yield strength of the eutectic alloy was determined using a MTS 810 mechanical testing system. The two phase alloy was subjected to mechanical testing at temperatures shown in FIG. 3. The yield strength at 294 K (room temperature) was determined to be about 750 MPa. High strength was maintained to about 675 K (400° C.) after which the yield strength decreased but was still approximately 200 MPa at 975 K (700° C.).

The stress versus strain curve of FIG. 4 shows results of a tensile test performed with a MTS 810 mechanical testing system at room temperature. A 1.75 inch sample of Fe₃₀Ni₂₀Mn₃₅Al₁₅ was deformed about 0.35 inches prior to fracturing in a ductile manner under an applied stress of about 1200 MPa. The observed percent elongation was about 20% at room temperature.

As shown in FIG. 5, the tensile and yield strengths of Fe₃₀Ni₂₀Mn₃₅Al₁₅ are significantly greater than those of many known alloys, including austenitic, ferritic and mild steel. The ductility, measured as percent elongation, is also comparable to ferritic and mild steel, as shown in FIG. 6.

FIGS. 7A-B illustrate the microstructure of an alloy with the nominal composition Fe₃₀Ni₂₀Mn₃₅Al₁₅, prepared by arc melting a mixture of Fe, Ni, Mn and Al in a water-cooled copper crucible under an argon atmosphere (e.g. 24.3 g Fe, 17 g Ni, 27.8 g Mn and 5.9 g Al). The ingot was melted for 1 min and flipped using a probe inside the chamber. The process was repeated three times to ensure homogeneity then the ingot was cooled rapidly to room temperature. FIGS. 7A and 7B show a scanning electron microscope (SEM) images of Fe₃₀Ni₂₀Mn₃₅Al₁₅ alloys; FIG. 7B is at higher magnification than FIG. 7A. The dark and light regions in FIGS. 7A and 7B are lamellae of f.c.c. and B2 phases, respectively. In FIG. 7A, some of the lamellae were normal to the surface being examined such that only the cross-section of the lamellae was seen.

In contrast, FIG. 8 is a TEM micrograph of a specimen quenched from 1623 K directly into water and then annealed at 473 K for 1 h to reduce residual stresses. The B2 phase spacing was measured as 50-70 nm and the f.c.c. phase spacing was measured as 80-150 nm. In FIG. 8, some of the B2 and f.c.c. elongated structures are sufficiently curved that they enclose each other when viewed in a polished cross-section.

Example 2 Preparation and Characterization of Fe_(x)Ni_(50-x)Mn_(50-y)Al_(y)±5%

Various alloys are cast with a composition:

Fe_(x)Ni_(50-x)Mn_(50-y)Al_(y),  Formula (2)

where x ranges from 25 to 35 atomic percent plus or minus 5%, and y ranges from 10 to 20 atomic percent plus or minus 5%.

The alloys are cast using the aforementioned arc melting technique and heated to a eutectic transformation temperature range of between about 1210-1290° C., or between about 1212-1250° C. or between about 1214-1230° C. The alloys are expected to be strong and ductile with a range of mechanical properties that can be manipulated by composition variations within the disclosed range.

FIGS. 9A-C show SEM images of (a) Fe₂₈Ni₁₈Mn₃₃Al₂₁; (b) Fe₂₉Ni₁₉Mn₃₄Al₁₈ and (c) Fe₂₈Ni₂₁Mn₃₃Al_(18.)In these images the B2 phase is brighter and the f.c.c. phase darker. The B2 and f.c.c. structures of the Fe₂₈Ni₁₈Mn₃₃Al₂₁ alloy are finer than those of the other two alloys.

FIGS. 10A-D show SEM images of: (a) Fe₃₁Ni₁₈Mn₃₈Al₁₃; (b) Fe₂₉Ni₁₉Mn₃₈Al₁₄; (c) Fe₃₃Ni₁₉Mn₃₄Al₁₄ and (d) Fe₃₆Ni₁₈Mn₃₃Al_(13.)

FIG. 11 shows stress-strain curves of several FeNiMnAl alloys: Fe₃₁Ni₁₈Mn₃₈Al₁₃; Fe₂₉Ni₁₉Mn₃₈Al₁₄; Fe₃₃Ni₁₉Mn₃₄Al₁₄ and Fe₃₆Ni₁₈Mn₃₃Al₁₃ and Fe₃₀Ni₂₀Mn₃₅Al_(18.)

Table 1 gives the percent elongation to failure, yield stress and ultimate tensile strength measured for alloys with less than 15 at % aluminum. (strain rate 5×10⁻⁴ s¹). Table 2 gives hardness measurements for alloys with less than 15 at % aluminum.

TABLE 1 Composition elongation/% yield stress/MPa UTS/MPa Fe₃₁Ni₁₈Mn₃₈Al₁₃ 21.6 460 780 Fe₂₉Ni₁₉Mn₃₈Al₁₄ 19.4 589 891 Fe₃₃Ni₁₉Mn₃₄Al₁₄ 15.7 614 916 Fe₃₆Ni₁₈Mn₃₃Al₁₃ 26.1 463 814 Fe₃₀Ni₂₀Mn₃₅Al₁₅ 6.5 820 1053

TABLE 2 Vickers hardness/ Composition vpn Fe₂₉Ni₁₉Mn₃₄Al₁₈ 375 ± 9 Fe₂₈Ni₁₈Mn₃₃Al₂₁  467 ± 18 Fe₂₈Ni₂₁Mn₃₃Al₁₈ 343 ± 6 Fe₃₃Ni₁₉Mn₃₄Al₁₄ 240 ± 5 Fe₃₆Ni₁₈Mn₃₃Al₁₃ 208 ± 7 Fe₂₉Ni₁₉Mn₃₈Al₁₄ 247 ± 6 Fe₃₁Ni₁₈Mn₃₈Al₁₃ 213 ± 5

Example 3

Characterization of a Phase Diagram Near a Eutectic Transformation

A portion of a phase diagram near a eutectic transformation may be constructed by varying percentages of Fe, Ni, Mn, Al and M as described in the context of Formula (1), except the subscripts a, b, c, d, and e, may be any value. The constituents are processed as described in Examples 1 and 2 to ascertain the presence or absence of eutectic transformation products. The preferred metals include combinations of Fe, Ni, Mn, and Al, in which case the ranges for x and y shown in Formula (2) may be any value. When adjusting the respective subscripts a, b, c, d, e, x and/or y, it is suggested to increase or decrease the individual ranges or combinations of ranges in steps of five percent from the values shown in Formulas (1) and (2), at least until the resulting alloy does not show evidence of a eutectic transformation. For alloys that contain four or five constituents, it is routine in the art that several hundred castings are needed to fully characterize the phase diagram around a eutectic transformation.

Example 4 Preparation and Characterization of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with Cr Additions

Fe₃₀Ni₂₀Mn₃₅Al₁₅ with different additions of Cr (0-8 at. %) were prepared by arc melting a mixture of elemental Cr, Fe, Ni, Mn and Al with purities >99.8% in a water-cooled copper crucible under an argon atmosphere. The ingot was melted and flipped three times using a probe inside the chamber to ensure homogeneity, followed by cooling rapidly to room temperature.

X-ray diffraction (XRD) was performed on the alloys using a Rigaku D/Max 2000 diffractometer with Cu Kα radiation operated at 40 kV and 300 mA. Measurements were performed by step scanning 2θ from 20° to 120° with a 0.02° step size. A count time of 1 s per step was used. Specimens, cut from the ingot, were polished using increasingly fine grade of silicon carbide papers followed by 0.3 μm alumina powder to a mirror finish. The surface was etched using 4% nitric acid for ˜5 s followed by rinsing in water. Specimens were examined in a FEI XL-30 field emission gun scanning electron microscope (SEM) operated at 15 kV. Discs of 3 mm in diameter and 100 μm thick, from either before or after ˜5% strain under compression, were electropolished using 30% nitric acid in methanol at 253 K in a Struers Tenupol 5, and washed alternatively in ethanol and methanol for three cycles followed by a final rinse in methanol. The typical voltage and current were 11 V and 100 mA, respectively. The resulting thin foils were examined using an FEI Tecnai F20 field emission gun transmission electron microscope (TEM) equipped with energy dispersive X-ray spectrometry (EDS), operated at 200 kV.

Flat tensile specimens of dog bone geometry with a gauge length of 10 mm and thickness of ˜1.27 mm were produced for tensile tests. The specimens were polished using silicon carbide papers and finished with 0.3 μm alumina powder to reduce surface defects. Tensile tests were performed at different strain rates (5×10⁻⁶ s⁻¹−5×10⁻¹ s⁻¹) at room temperature and at a strain rate of 5×10⁻⁴ s⁻¹ at temperatures from 300 K to 1000 K.

Secondary electron (SE) images of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with different Cr additions are shown in FIGS. 12A-D. For Cr additions of 2 at. %, 4 at. % and 6 at. %, the material exhibited a lamellae microstructure with lamellar spacing of the order of a few hundred nanometers. Previously, the microstructure of the as-cast Cr-free alloy was reported to consist of 200 nm wide B2 lamellae and 500 nm wide f.c.c. lamellae (Liao and Baker, Mater. Charact. 59 (2008) 1546-1549). However, the addition of 8 at. % Cr produced a finer more complicated microstructure, as shown in FIG. 12D, with some cuboidal phases present. In the images of FIGS. 12A-12D, the B2 lamellae are brighter contrast than the f.c.c. lamellae.

XRD patterns of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr and 8 at. % Cr, see FIG. 13A-B, show only the presence of B2 and f.c.c. phases, i.e. the same as those in as-cast Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅ (Liao and Baker, 2008). The reason for the microstructure change when 8 at. % Cr is present is not clear.

FIG. 14A is a bright field (BF) TEM image of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 4 at. % Cr, showing the lamellar structure. In the image, the f.c.c. phase is brighter and the B2 phase darker. FIGS. 14B-C confirms that the two phases present are B2 and f.c.c. FIGS. 14D-E shows Xray spectra from the two phases, and Table 3 lists the measured compositions of the phases. The latter demonstrate that the Cr partitions largely into the f.c.c. phase.

TABLE 3 F.C.C. B.C.C Fe 36.7 5.3 Ni 12.5 39.0 Mn 34.9 14.7 Al 11.8 40.6 Cr 4.2 0.6

FIG. 15A is a BF TEM image of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 8 at. % Cr showing an overview of the microstructure. In this figure, the B2 phase is dark and the f.c.c. phase light. The structure is finer than that of FIGS. 12A to 12C, with an average spacing between B2 structures of ˜100 nm. The structure also includes cuboidal particles. FIG. 15B shows the cuboidal phase in more detail with an embedded selected area diffraction (SAD) pattern indicating that the cuboidal phases are again B2.

The effects of Cr on the room temperature strength and

ductility of Fe₃₀Ni₂₀Mn₃₅Al₁₅ are illustrated in FIGS. 16A-B. Tensile tests, performed at a strain rate of 5×10⁻⁴ s⁻¹, showed that the ductility increased steadily with increasing Cr up to 6 at. %, but decreased dramatically at 8 at. % Cr addition. The latter feature reflects the change in microstructure at 8 at. % Cr. The elongation to fracture measured from specimens increased from 6.5% when no Cr was present to 17.8% when 6 at. % Cr was added, but decreased dramatically to only 6.3% when at 8 at. % Cr. Concurrent with the increase in ductility, the yield stress decreased from 820 MPa for the Cr-free alloy to 679 MPa for the alloy containing 6 at. % Cr, and then increased to 819 MPa when the Cr addition was increased to 8 at. %, see FIG. 16B. The ultimate tensile stress decreased monotonically as the Cr content increased.

BF TEM analysis of Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr after ˜5% strain showed dislocations were generated within f.c.c. lamellae and piled up at interphase interfaces. However, no dislocations were observed in the B2 phase. It is evident that the deformation mainly occurred in f.c.c phase, while the B2 phase acted as obstacles to moving dislocations. Room temperature tensile tests were also performed as a function of strain rate for Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr, see FIG. 17A. The yield stress, ultimate tensile stress (UTS) and elongation to fracture are essentially insensitive to the strain rate as shown in FIG. 17B. This is in sharp contrast to Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅, which exhibited a strong strain rate dependence of the ductility and, hence, UTS at room temperature (Liao et al., 2011, Intermetallics, 19, 1533-1537). Liao et al. showed that the elongation to fracture and ultimate tensile stress of Fe₃₀Ni₂₀Mn₃₅Al₁₅ increased with increasing strain rate below 3×10⁻³ s⁻¹, but were independent of stain rate at ˜10.5% and 840 MPa for strain rates ≧3×10⁻³ s⁻¹. The elongation was only 0.7% at a strain rate of 3×10⁻⁶ s⁻¹ (Lio et al., 2011). Thus, it is clear that the Cr addition suppressed the environmental embrittlement that is observed in Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅ at low strain rate.

Fracture surfaces from the 6 at. % Cr-modified Fe₃₀Ni₂₀Mn₃₅Al₁₅

tensile tested at a strain rate of 5×10⁻⁶ s⁻¹ at room temperature showed ductile tearing with elongated dimples (See FIGS. 18A-B). By comparison the fracture surfaces from the Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅ alloy tested under the same testing conditions, see FIGS. 18C-D, show mixed mode fracture, i.e. both dimple type ductile fracture (FIG. 18C) and brittle transgranular cleavage (FIG. 18D). In other words, the fracture mode at low strain rate changes when chromium is present. FIG. 19A shows strain-stress curves for Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr for tensile tests performed at a variety of temperatures up to 1000 K at a strain rate of 5×10⁻⁴ s⁻¹ The yield stress and ultimate tensile stress as a function of temperature are plotted in FIG. 19B; the yield stress of Cr-free Fe₃₀Ni₂₀Mn₃₅Al₁₅ is shown for comparison (Liao and Baker, 2011, J. Mater. Sci. 46, 2009-2017). The yield stress and ultimate tensile stress of as cast Fe₃₀Ni₂₀Mn₃₅Al₁₅ with 6 at. % Cr slightly decreased from 300 to 600 K, followed by a sharp drop after 600 K. The behavior is the same as the Cr-free alloy. The yield stress is only ˜150 MPa at 1000 K, i.e. again very similar to the value for the Cr-free alloy. At the highest temperature the yield strength and UTS have similar values due to the lack of work-hardening. The elongation to failure was largely independent of temperature up to 800 K with a small maximum at the latter temperature, but decreased slightly at 900 K and 1000 K, see FIG. 20. The reduction in area at the neck is also plotted in FIG. 20 as a function of temperature and shows a somewhat similar trend to the elongation as a function of temperature, with a strong maximum at 800 K.

The fracture surfaces of the 6 at. % Cr-modified Fe₃₀Ni₂₀Mn₃₅Al₁₅ tensile tested at elevated temperature at strain rate of 5×10⁻⁴ s⁻¹ were also examined. Examination of typical fracture surfaces at 500 K and 800 K showed ductile, dimple-type rupture mode was found at both temperatures, but the size and depth of the dimples was greater at 800 K compared to those at 500 K, which is to be expected since the tests at 800 K showed maxima in both elongation and reduction in area.

Chromium additions up to 6 at. % have three intriguing effects on the room temperature mechanical behavior of Fe₃₀Ni₂₀Mn₃₅Al₁₅, while producing little change in the microstructure.

First, Cr decreases the yield strength. Since plastic deformation at room temperature is accommodated solely by plastic deformation of the f.c.c. phase and as shown by EDS, most of the Cr partitions into the f.c.c. phase (FIGS. 14D-E)., it is evident that the effect of Cr is to soften this phase.

The second intriguing effect is the increase in ductility with increasing Cr (up to 6 at. %). Previous work showed that dislocation pile-ups in the f.c.c. phase at the f.c.c./B2 interface produced cracking in the B2 phase (Liao and Baker, 2011, Mater. Sci. Eng.: A 528, 3998-4008), which ultimately led to failure. Without wishing to be bound by any particular belief, a reduction in the length of dislocation pile-ups at a given stress due to Cr additions may delay this cracking until larger strains are reached. The third effect is that the Cr addition (particularly for 6 at. % Cr) suppresses the environmental embrittlement problem observed in unalloyed Fe₃₀Ni₂₀Mn₃₅Al₁₅ when tested at slow strain rates at room temperature (Liao et al., 2011). Without wishing to be found by any particular believe, there may be several explanations for the effect of chromium on the environmental embrittlement in Fe₃₀Ni₂₀Mn₃₅Al₁₅. First, as noted above, the Cr addition may increase the stacking fault energy and lead to wavy slip, as observed by McKamey et al. in Fe₃Al (J. Mater. Res. 1989, 4, 1156-1163; Scr. Metall, 1988, 22, 1670-1681). This would reduce dislocation pile-ups, and, hence, stress concentrations. It may also reduce the ability of gliding dislocations to transport atomic hydrogen, produced by the reaction between water vapor and Al, into the material. However, it is worth noting that Schramm and Reed (1975, Metall. Trans. A, 6, 1345-1351) suggested that Cr increased the stacking fault energy of austenitic stainless steels.

Second, the Cr addition may result in a change in oxide chemistry and properties, or a change in the kinetics of oxide formation (Mckamey and Liu, 1990, Scr. Metall. Mater. 24, 2119-2122). It is likely that mixed oxides are present on the surface when Cr is present. Both of these mechanisms will minimize the environmental embrittlement by reducing the water-vapor reaction. Finally, the possible mechanisms responsible for the beneficial effect of Cr addition on the embrittlement of Ni₃(Si, Ti) alloys suggested by Ma et al. (1995, Scr. Metall. Mater. 32, 1025-1029), i.e. blocking site occupation and/or the diffusion of hydrogen along grain boundaries; affecting both the water decomposition process and the subsequent hydrogen absorption processes; and enhancing the cohesive strength of the grain boundaries, may also be important in the alloy tested here.

The temperature dependence of the yield strength of the 6 at. % Cr modified alloy is essentially the same as the Cr-free alloy (Liao and Baker, 2011, J. Mater. Sci. 246, 2009-2017). Differential thermal analysis showed that the melting point of 6 at. % Cr modified Fe₃₀Ni₂₀Mn₃₅Al₁₅ was ˜1500 K. The yield strength of Cr modified Fe₃₀Ni₂₀Mn₃₅Al₁₅ is only weakly dependent on temperature up to 600 K, which is ˜0.45 Tm. The rapid decrease in yield stress with increasing temperature above 600 K has been observed in many B2 compounds above 0.45 Tm and is associated with the onset of diffusive processes (Baker, 1995, Mat. Sci. Eng. A-Struc. Mater. Prop. Microstruct. Process. 192, 1-13). The B2 phase in Fe₃₀Ni₂₀Mn₃₅Al₁₅ experiences a brittle-to-ductile transition with increasing temperature, as in many B2 compounds (Baker, 1995). At lower temperatures, the B2 phase is hard and acts as an obstacle to dislocation motion, eventually fracturing due to stress concentrations from dislocation pile-ups in the f.c.c. phase (Liao and Baker, 2011, Mater. Sci. Eng. A 528, 3998-4008). In contrast, at higher temperatures the B2 lamellae undergoes plastic deformation, and the overall yield stress the elongation to failure and reduction in area at 800 K followed by a decrease to their lowest values at 900 K and 1000 K. Dimple-type ductile fracture was observed at all temperatures. In contrast, B2 compounds typically show increasing ductility with increasing temperature (Baker 1995) For example, for boron-doped (500 ppm) B2 Fe-45Al the elongation to failure increased from _(—)3% at room temperature with increasing temperature up to 40% at 800 K, but this was followed by a sharp drop to 20% above that temperature when necking occurred prior to failure (Klein and Baker, 1994, Scr. Metall. Mater., 30, 1413-1417). However, in contrast to the present work, the reduction in area increased continuously with increasing temperature reaching a value of 90% at 1000 K (Klein and Baker, 1994). The reason for the decrease in elongation at higher temperatures is because the FeAl no longer showed any work-hardening. Thus, once a neck formed during deformation, there was no capacity for work hardening to offset the reduction in area at the neck.

The lack of a significant increase in ductility up to 700 K in Fe₃₀Ni₂Mn₃₅Al₁₅ may reflect the fact that the B2 phase does not undergo plastic deformation up to that temperature as indicated by the slight change in yield strength with increasing temperature up to that point. The peak at 800 K is presumably related to the fact that the B2 phase now undergoes plastic deformation as reflected in a substantial decrease in yield strength at that temperature. Similar to FeAl, the reduction in elongation observed at 900 K and 1000 K is due to the lack of work-hardening at these temperatures, which produces unstable necking. However, it is somewhat surprising that, unlike FeAl (Klein and Baker, 1994), the reduction in area also decreases at 900 K and 1000 K. This probably reflects the extreme cavitation in the neck.

Additional details may be found in Meng et al., 2013, Mater. Sci. Engr. A, 586, 45-52, hereby incorporated by reference.

It is understood for purposes of this disclosure, that various changes and modifications may be made to the disclosed embodiments that are well within the scope of the present compositions and methods. Numerous changes may be made which will readily suggest themselves to those skilled in the art and which are encompassed in the spirit of the compositions and methods disclosed herein and as defined in the appended claims.

All references throughout this application, for example patent documents including issued or granted patents or equivalents; patent application publications; and non-patent literature documents or other source material; are hereby incorporated by reference herein in their entireties, as though individually incorporated by reference, to the extent each reference is at least partially not inconsistent with the disclosure in this application (for example, a reference that is partially inconsistent is incorporated by reference except for the partially inconsistent portion of the reference).

All patents and publications mentioned in the specification are indicative of the levels of skill of those skilled in the art to which the invention pertains. References cited herein are incorporated by reference herein in their entirety to indicate the state of the art, in some cases as of their filing date, and it is intended that this information can be employed herein, if needed, to exclude (for example, to disclaim) specific embodiments that are in the prior art. For example, when a compound is claimed, it should be understood that compounds known in the prior art, including certain compounds disclosed in the references disclosed herein (particularly in referenced patent documents), are not intended to be included in the claim

As used herein, “comprising” is synonymous with “including,” “containing,” or “characterized by,” and is inclusive or open-ended and does not exclude additional, unrecited elements or method steps. As used herein, “consisting of” excludes any element, step, or ingredient not specified in the claim element. As used herein, “consisting essentially of” does not exclude materials or steps that do not materially affect the basic and novel characteristics of the claim. Any recitation herein of the term “comprising”, particularly in a description of components of a composition or in a description of elements of a device, is understood to encompass those compositions and methods consisting essentially of and consisting of the recited components or elements. The invention illustratively described herein suitably may be practiced in the absence of any element or elements, limitation or limitations which is not specifically disclosed herein. 

What is claimed is:
 1. An alloy comprising lamellae of a B2 ordered phase and lamellae of a face-centered cubic (f.c.c.) phase, wherein the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum, and 2.5% to 5% chromium, wherein the composition is described in terms of atomic percentages.
 2. The alloy of claim 1 wherein lamellae of the B2 ordered phase alternate with lamellae of the f.c.c. phase.
 3. The alloy of claim 1, wherein the average composition of the alloy comprises from 3% to 5% chromium.
 4. The alloy of claim 1, wherein the concentration of nickel and aluminum in the B2 phase is greater than the amount or iron or manganese and the concentration of iron and manganese in the f.c.c. phase is greater than the amount of nickel or aluminum.
 5. The alloy of claim 1, wherein the average thickness of the lamellae of the B2 phase is from 150 nm to 750 nm.
 6. The alloy of claim 1, wherein the average thickness of the lamellae of the B2 phase is from 150 nm to 500 nm.
 7. The alloy of claim 1, wherein the average thickness of the lamellae of the f.c.c. phase is from 200 nm to 2000 nm.
 8. The alloy of claim 1, wherein the average thickness of the lamellae of the f.c.c. phase is from 300 nm to 1000 nm.
 9. The alloy of claim 1, wherein the elongation to fracture of the chromium-containing alloy is from 10% to 15% as measured in air at room temperature at a strain rate of 5×10⁻⁴ s⁻¹.
 10. An alloy comprising lamellae of a B2 ordered phase and lamellae of an f.c.c. phase, wherein the average composition of the alloy comprises from 27% to 33% iron, 17% to 23% nickel, 32% to 38% manganese and 12% to 18% aluminum, and greater than 5% to 7.5% chromium, wherein the composition is described in terms of atomic percentages.
 11. The alloy of claim 10, wherein lamellae of the B2 ordered phase alternate with lamellae of the f.c.c. phase.
 12. The alloy of claim 10, wherein the average composition of the alloy comprises 5.5% to 6.5% chromium.
 13. The alloy of claim 10, wherein the average composition of the alloy comprises 6% chromium.
 14. The alloy of claim 10, wherein the concentration of nickel and aluminum in the B2 phase is greater than the amount or iron or manganese and the concentration of iron and manganese in the f.c.c. phase is greater than the amount of nickel or aluminum.
 15. The alloy of claim 10, wherein the average thickness of the lamellae of the B2 phase is from 150 nm to 750 nm.
 16. The alloy of claim 10, wherein the average thickness of the lamellae of the B2 phase is from 150 nm to 500 nm.
 17. The alloy of claim 10, wherein the average thickness of the lamellae of the f.c.c. phase is from 200 nm to 2000 nm.
 18. The alloy of claim 10, wherein the average thickness of the lamellae of the f.c.c. phase is from 300 nm to 1000 nm.
 19. The alloy of claim 10 wherein the elongation to fracture of the chromium-containing alloy is from 15% to 25% as measured in air at room temperature at a strain rate of 5×10⁻⁴ s⁻¹. 